Valve terminology varies according to the industry (e.g., pipeline or oil field service) in which the valve is used. In some applications, the term “valve” means just the moving element, whereas in other applications, the term “valve” includes the moving element, the valve seat, and the housing that contains the moving element and the valve seat. A valve assembly of the present invention comprises a valve body (the moving element) and a corresponding valve seat, the valve body typically incorporating an elastomeric seal within a peripheral seal retention groove.
Valve assemblies of the present invention can be mounted in the fluid end of a high-pressure pump incorporating positive displacement pistons or plungers in multiple cylinders. Such a valve assembly may incorporate a web-seat, stem-guided valve design or a full open seat design, since either design can be adapted for high pressures and repetitive impact loading of the valve body and valve seat. These severe operating conditions have in the past often resulted in leakage and/or premature valve failure due to metal wear and fatigue. In overcoming such failure modes, special attention is focused in the present invention on valve sealing surfaces where the valve body contacts the valve seat intermittently for reversibly blocking fluid flow through a valve assembly.
Valve assembly sealing surfaces are subject to exceptionally harsh conditions in exploring and drilling for oil and gas, as well as in their production. For example, producers often must resort to “enhanced recovery” methods to insure that an oil well is producing at a rate that is profitable. And one of the most common methods of enhancing recovery from an oil well is known as fracturing. During fracturing, cracks are created in the rock of an oil bearing formation by application of high hydraulic pressure. Immediately following fracturing, a slurry comprising sand and/or other particulate material is pumped into the cracks under high pressure so they will remain propped open after hydraulic pressure is released from the well. With the cracks thus held open, the flow of oil through the rock formation toward the well is usually increased.
The industry term for particulate material in the slurry used to prop open the cracks created by fracturing is the propend. And in cases of very high pressures within a rock formation, the propend may comprise extremely small aluminum oxide spheres instead of sand. Aluminum oxide spheres may be preferred because their spherical shape gives them higher compressive strength than angular sand grains. Such high compressive strength is needed to withstand pressures tending to close cracks that were opened by fracturing. Unfortunately, both sand and aluminum oxide slurries are very abrasive, typically causing rapid wear of many component parts in the positive displacement plunger pumps through which they flow. Accelerated wear is particularly noticeable in plunger seals and in the suction (i.e., intake) and discharge valve assemblies of these pumps.
A valve assembly 10 (comprising a valve body 20 and valve seat 40) representative of a typical full open design valve and seat for a fracturing plunger pump is schematically illustrated in FIG. 1. FIG. 2 shows how sand and/or aluminum oxide spheres may become trapped between sealing surface 21 of valve body 20 and sealing surface 41 of valve seat 40 as the suction valve assembly 10 closes during the pump's pressure stroke.
The valve assembly 10 of FIG. 1 is shown in the open position. FIG. 2 shows how accelerated wear begins shortly after the valve starts to close due to back pressure. For valve assembly 10, back pressure tends to close the valve when downstream pressure exceeds upstream pressure. For example, when valve assembly 10 is used as a suction valve, back pressure is present on the valve during the pump plunger's pressure stroke (i.e., when internal pump pressure becomes higher than the pressure of the intake slurry stream. During each pressure stroke, when the intake slurry stream is thus blocked by a closed suction valve, internal pump pressure rises and slurry is discharged from the pump through a discharge valve. For a discharge valve, back pressure tending to close the valve arises whenever downstream pressure in the slurry stream (which remains relatively high) becomes greater than internal pump pressure (which is briefly reduced each time the pump plunger is withdrawn as more slurry is sucked into the pump through the open suction valve).
When back pressure begins to act on a valve, slurry particles become trapped in the narrow space that still separates the sealing metal surfaces of the valve body and seat. This trapping occurs because the valve is not fully closed, but the valve body's elastomeric seal has already formed an initial seal against the valve seat. The narrow space shown in FIG. 2 between metallic sealing surfaces 21 and 41 of the valve body and valve seat respectively is typically about 0.040 to about 0.080 inches wide; this width (being measured perpendicular to the sealing surfaces of the valve body and seat) is called the standoff distance. The size of the standoff distance is determined by the portion of the valve body's elastomeric seal that protrudes beyond the adjacent valve body sealing surfaces to initially contact, and form a seal against, the valve seat. As schematically illustrated in FIG. 2, establishment of this initial seal by an elastomeric member creates a circular recess or pocket that tends to trap particulate matter in the slurry flowing through the valve.
Formation of an initial seal at a valve that is closing under back pressure immediately stops slurry flow through the valve. Swiftly rising back pressure tends to drive slurry backwards through the now-sealed valve, but since back-flow is blocked by the initial valve sealing, pressure builds rapidly on the entire valve body. This pressure acts on the area of the valve body circumscribed by its elastomeric seal to create a large force component tending to completely close the valve. For example, a 5-inch valve exposed to a back pressure of 15,000 pounds per square inch will experience a valve closure force that may exceed 200,000 pounds.
The large valve closure force almost instantaneously drives the affected valve assembly, whether suction or discharge, to the fully closed position where the metal sealing surface of the valve body contacts the corresponding metal sealing surface of the valve seat. As the valve body moves quickly through the standoff distance toward closure with the valve seat, the elastomeric seal insert is compressed, thus forming an even stronger seal around any slurry particles that may have been trapped between the seal insert and the valve seat.
Simultaneously, the large valve closure force acting through the standoff distance generates tremendous impact energy that is released against the slurry particles trapped between the metallic sealing surfaces of the valve body and the valve seat. As shown in FIG. 3, the slurry particles that are trapped between approaching valve sealing surfaces 21 and 41 are crushed.
In addition to the crushing action described above, slurry particles are also dragged between the valve sealing surfaces in a grinding motion. This grinding action occurs because valve bodies and seats are built with complementary tapers on the sealing surfaces to give the valve assembly a self-alignment feature as the valve body closes against the seat. As the large valve closing force pushes the valve body into closer contact with the seat, the valve body tends to slide down the sealing surface taper by a very small amount. Any crushed slurry particles previously trapped between the sealing surfaces are then ground against these surfaces, resulting in extreme abrasive action.
To limit sealing surface erosion due to this abrasion, valve bodies and seats have in the past been heat-treated to harden and strengthen them. Typical heat treatment methods have included carburizing, as well as hardening by induction heating and flame hardening. All of these hardening processes depend on quenching (i.e., rapid cooling) of the valve components after they have been uniformly heated, preferably slightly above a critical temperature (called the upper transformation temperature).
When a steel object is uniformly heated to a temperature slightly above its upper transformation temperature, all of the steel in the object assumes a face-centered cubic crystal lattice structure known as austenite. When the object is quenched below this temperature, other crystal lattice structures are possible. If quenched uniformly, the other crystal lattice structures tend to appear uniformly throughout the object. But if certain portions of the object are cooled at rates different from those applicable to other portions of the object, then the crystal lattice structure of the cooled object may be non-uniform.
Further, if steel is heated too far above its upper transformation temperature before quenching, its grain structure may be unnecessarily coarsened, meaning that the steel will then be less tough and more brittle after quenching than it would have been if its maximum temperature had been closer to its upper transformation temperature. It is therefore important that heat treatments for a particular steel be applied uniformly when uniform results are desired, and it is further important that maximum temperatures not be so high as to adversely affect the steel's grain structure.
Maximum heat-treatment temperatures for different steels vary because they are closely related to each steel's upper transformation temperature (which depends on that particular steel's composition). For example, carbon steel may have an upper transformation temperature as low as about 1333 degrees F., whereas high-alloy steels may have upper transformation temperatures of more than 2000 degrees F. The upper transformation temperature of the steel traditionally used in manufacturing high-pressure valve assemblies is about 1650 degrees F.
As an example of changes that can occur in a steel's crystal lattice structure at temperatures around its upper transformation temperature, consider a low-carbon steel. Such steels, commonly comprising iron and about 0.2% carbon with small amounts of alloying elements, are often used when carburizing is desired for hardening. As this steel begins cooling from a temperature slightly above its upper transformation temperature, its crystal lattice is 100% face-centered cubic (i.e., austenite), but as it cools the steel begins to assume other crystal lattice structures (typically referred to as martensite). Body-centered cubic forms are favored by relatively slow cooling, whereas body-centered tetragonal forms are favored by faster cooling. As cooling progresses, the percentage of austenite tends to decrease. And by the time the steel cools to a temperature of about 1333 degrees F. (called the lower transformation temperature), most of the austenite has been transformed to one or more other crystal lattice forms. Hence, at temperatures below about 1333 degrees F., little or no austenite exists and there are no further significant changes in the relative percentages of other crystal lattice forms present.
The above example of progressive changes in a low-carbon steel's crystal lattice differs considerably from the changes that would occur in higher alloy types of steel. Certain stainless steels, for example, can retain an austenitic lattice structure even at room temperature. In particular, the presence of nickel in steel alloys is observed to be associated with retention of austenitic lattice structure at temperatures below the lower transformation temperature.
The ability to predict the relative percentages of different crystal lattice forms present at different stages of a heat treatment allows adjustment of a wide variety of a steel's important physical properties to adapt it for specialized applications. An example of such an adaptation process comprises heating a steel to a predetermined temperature to within or slightly above a particular range (called the transfer temperature range) between the steel's upper transformation temperature and about 1333 degrees F. Following such heating, the steel is cooled (quenched) according to protocols developed to enhance physical properties such as hardness.
Quenching is performed primarily to influence the formation of a desirable crystal lattice and/or grain structure in a cooled metal, a grain being a portion of the metal having external boundaries and a regular internal lattice. Quenching may be accomplished, for example, simply by immersion of a heated metal object in water or oil. Certain tool steels may even be quenched by gas, but the carbon steels traditionally used for valve seats can not be gas-quenched if they are to develop the hardness, strength and toughness necessary for use in high-pressure valves.
Transitions between crystal structures that occur in association with quenching are not instantaneous, so the rate of cooling may be adjusted to favor development of more desirable crystal and/or grain structures with their corresponding beneficial material properties (e.g., tensile strength, hardness, ductility, toughness, etc.). Further, quenching may optionally be followed by tempering, wherein metal is reheated to a temperature below its lower transformation temperature before finally returning to room temperature. Tempering is particularly useful with relatively high-strength alloy steels because it makes the steel tougher and more ductile (as in tempered martensite) by reducing internal stresses that would otherwise tend to make steel brittle and prone to cracking (e.g., untempered martensite).
Heat treating of metals has been extensively studied, and many desirable properties may be obtained in metals through elaborate quench and temper protocols that have been experimentally developed. But preferred heat treatments are highly specific to particular alloys, so there may be no single optimal heat treatment for a component such as a valve seat comprising, for example, a high-alloy sealing surface inlay on a carbon steel substrate. Indeed, even the most careful use of heat treatments to favor development of hard sealing surfaces on strong, tough substrates has not proven effective for extending the service life of valves traditionally used for high-pressure abrasive slurries. Thus, engineers have long sought better methods of hardening valve sealing surfaces at acceptable cost.
For example, incorporation of metallic carbides in sealing surfaces has been investigated because some metallic carbides are extremely hard and wear-resistant. But such carbides do not bond well with the low-carbon steels commonly used in high pressure valve seats. Hence, when metallic carbide inlays are applied to such valve seat substrates, they must actually be held in place by some type of binder which itself forms an adequate bond with the valve seat substrate steel.
To facilitate mixing metallic carbides with binders, the carbides are made commercially available in powder form. Such powders (e.g., carbides of vanadium, molybdenum, tungsten or chromium) are formed by casting the pure carbides and then crushing them into the desired particle size. A binder (comprising, e.g., cobalt, chromium, and/or nickel) is then added to the crushed carbide powders, but there is little or no opportunity for the binder to alloy with the carbides.
Metallic carbide particles thus bound as an inlay on a steel substrate are called cemented carbides, and they comprise a matrix consisting of a dispersion of very hard carbide particles in the (relatively softer) binder. The resulting cemented carbide inlays are thus not homogeneous, so they do not possess the uniform hardness that would ideally be desired for good abrasion resistance and toughness in valve sealing surfaces. One problem associated with this inhomogeneity becomes evident because the crushing and grinding of slurry particles between valve sealing surfaces during valve closure produces a variety of slurry particle sizes, some so fine that they are smaller than the spacing between the carbide particles in the cemented carbide inlay. These fine slurry particles are very abrasive, and they can fit between the carbide particles to rapidly wear away the relatively soft binder holding the carbide particles in place. Thus loosened (but not actually worn down), the carbide particles can simply be carried away by the slurry stream, leaving the remainder of the inlay binder exposed to further damage by the abrasive slurry.
Notwithstanding the above problems, cemented carbides, particularly those applied by gas-fueled or electrically-heated welding equipment, have been widely used to reduce abrasion damage in various industrial applications. But weld-applied carbide inlays have not been found acceptable in high pressure valves. Rather, the repetitive high-impact loading common in such valves, combined with cemented carbide's brittle nature and tendency to crack, has tended to cause premature (often catastrophic) valve failures. Thus, a long-felt need remains for better technology that can be economically applied to harden valve sealing surfaces without imparting excessive brittleness or likelihood of fatigue cracking. Any such new technology should address the problems of carbide inhomogeneity noted above, as well as other problems (e.g., as discussed below) commonly associated with the welding technology used for carbide application.
Among the other welding-related problems seen in applications of cemented carbides to a steel substrate are those resulting from poor weld preparation and from relatively high amounts of carbides in the weld inlay. Additionally, the surface of a newly-applied carbide weld inlay is rather rough. But a valve seat requires a smooth finish to reduce friction and abrasion in contacts with the valve body and its elastomeric seal. A smooth finish also tends to prevent stress risers from developing on the wear surface, so most seats are machined with a 63 RMS (root mean square) surface finish. Unfortunately, traditional cemented carbides cannot be softened for machining by annealing so they must be ground to achieve a smooth surface. Grinding a valve seat surface, in turn, requires special and expensive tooling and fixturing, and the grinding itself is slow, labor-intensive and expensive.
Further problems related to weld-applied cemented carbides result from the large amounts of heat concentrated during welding on a relatively small area of steel (such as a valve mating surface inlay) which itself is adjacent to a larger mass of steel (such as a valve seat substrate). Portions of the substrate adjacent to the small heated volume of the inlay act as heat sinks, meaning that their temperature increases more slowly than the temperature of the heated (cemented carbide) inlay. Because of the temperature gradient thus created between the cemented carbide inlay and the steel substrate, and the generally different coefficients of expansion of these two materials, high shear stresses can develop between inlay and substrate during both heating and cooling. And unfortunately, a residual level of stress tends to persist after cooling and to be concentrated in both the base material and the cemented carbide inlay.
One reason that residual stresses persist after cooling of a weld-applied cemented carbide inlay on a larger steel substrate is that the rate of cooling of the liquid weld puddle is strongly influenced by the heat-sink effect of the relatively large mass of substrate, rather than being easily controllable through a quench and temper protocol which is applied to the mass as a whole. This heat-sink influence is termed the mass-quench effect, and it causes residual stresses that predispose weld-applied cemented carbide mating surfaces to premature cracking. While such cracks are tolerated in certain applications where the cracks do not significantly affect the performance of the part, the same can not be said of high pressure valve assemblies. On the contrary, cyclic fatigue associated with the repeated large impact loads experienced by these valves magnifies the deleterious effects of cracks and residual stresses. Premature catastrophic failures of valve bodies and/or seats are a frequent result.
Thus, surface cracks and internal stresses, combined with the inhomogeneity of the inlays noted above, constitute significant disadvantages of weld-applied cemented carbide inlays. And there is a further problem related to the mass-quench effect after welding. It is the appearance after cooling of two distinct thin layers of substrate steel just under the cemented carbide inlay. These two layers have undesirable properties and they arise because when an inlay is applied by welding, there are two heat-affected zones that result from the welding. The first heat-affected zone is a layer of steel immediately adjacent to the weld inlay. Steel in this first zone is heated above its transformation temperature by molten metal in the weld puddle, and as the weld puddle cools, steel of the first zone layer is quenched by the adjacent “semi-cool” substrate (i.e., the mass-quench effect). But the quench is partial and incomplete because of heat retained in the solidifying weld puddle. As a result, some grains of steel are never transformed from austenite to martensite, and the retained austenite in this layer is undesirable because of its brittleness. Equally detrimental is the fact that even the martensite that is formed in this layer is itself untempered and brittle. Fortunately, the brittleness of the steel in this first zone layer can be mitigated by a simple tempering heat treatment cycle, during which existing martensite is tempered, and simultaneously at least some of the retained austenite will be converted to tempered (i.e., less brittle) martensite. Any retained austenite, of course, will continue to contribute to brittleness in this first zone layer.
Adjacent to the first zone steel layer described above, a second heat-affected zone develops between the first layer and the remainder of the steel substrate. In this second zone layer the coolness of the adjacent substrate prevents steel from reaching its transformation temperature. Nevertheless, the heat transmitted through the first zone layer from the molten weld pool is sufficient to temper the steel in this second zone layer. Tempering increases the steel's ductility, but it also adversely affects strength, toughness, and other physical properties which are necessary for a valve seat to withstand the high impact loads characteristic of high-pressure pumps.
Unfortunately, the physical properties of tempered steel in the second zone layer can only be improved by a full body (i.e., both inlay and substrate) quench and temper heat treatment. As noted above, such an operation brings its own set of problems due to generally significant differences in the coefficents of thermal expansion of the steel substrate and the cemented carbide weld inlay.
Further, the preferred quench and temper protocols for the two portions also differ, and they can not be applied individually. Rather, the mass-quench effect dominates the cooling process and the result is that weld-applied cemented carbide inlays are not only inhomogeneous, but they are also unavoidably associated with undesirably high levels of residual internal stresses. These stresses aggravate the above noted tendency of such cemented carbide inlays to form cracks that can lead to premature catastrophic valve failures.
To address the problem of inhomogeneity in high-pressure valve seats, the combination of cemented carbide inlays on low-carbon steel substrates has been replaced experimentally with wear-resistant (relatively high-carbide) tool steel cladding on low carbon steel substrates. The tool steel cladding is commercially available as a powder in which all the elements have been mixed, melted and then gas atomized into spheres. High grades of these tool steel cladding powders are called P/M (for particle metallurgy) grades, and they generally cost at least 10 times per unit weight more than lower grade tool steels. Notwithstanding the high grade and high cost of the tool steel cladding however, these experimental valve seats have not been successful because the reheat treatment required to reduce the cladding's brittleness does not simultaneously cause development of the required strength and toughness in the low-carbon steel substrate.
The above-noted difficulty of reducing the brittleness of a relatively high carbide P/M inlay while simultaneously developing strength and toughness in a low-carbon steel substrate is addressed by substituting low grade tool steel (e.g., H13) for the low-carbon steel of the substrate. Residual internal stress is thereby reduced because a cladding matrix of high alloy P/M powder has a coefficient of thermal expansion which closely matches that of a low grade tool steel substrate. Such close matching of thermal expansion coefficients is not seen with inlays of either cemented carbide or tool steel on a low-carbon steel substrate. Further, during the melting and atomization of P/M alloys, the elements combine to form very fine carbides. Some of the carbon and other elements alloy with the iron to form very high alloy steel, and some of the carbides are then able to alloy with the steel. The combination of the high alloy steel and the very fine alloyed carbides give cladding comprising such P/M tool steel the effect of being of uniform hardness and homogeneity throughout.
For example, the high alloy P/M tool steel grades such as REX 121 (available from Crucible Materials) and Maxamet (available from Carpenter Steel) have coefficients of expansion very near the coefficient of expansion of a lower grade tool steel such as H13. Thus, a valve seat comprising an H13 substrate with an inlayed wear surface of REX 121 will develop very little residual stress when heat treated. The published composition of REX 121 metallic powder is 3.4% carbon, 4.0% chromium, 10.0% tungsten, 5.0% molybdenum, 9.5% vanadium, 9.0% cobalt, and the balance iron. The published composition of Maxamet metallic powder is 2.15% carbon, 4.75% chromium, 13.0% tungsten, 6.0% vanadium, 10.0% cobalt, and the balance iron. And the composition of H 13 is as follows: 0.4% Carbon, 0.35% Manganese, 1% Silicon, 5.2% Chromium, 1.3% Molybdenum, and 0.95% Vanadium; with the balance iron.
Note that the above P/M grades are heavy in carbide-forming alloys, as well as having elevated levels of carbon to form the carbides. Further, in addition to demonstrating uniform hardness, these substantially homogeneous tool steels have high strength and good toughness compared with cemented carbides. P/M grade tool steels are heat treatable and can be annealed to allow for easy machining before being rehardened and tempered.
The P/M grade tool steels also have very uniform and fine grain size, which eliminates segregation. Segregation is a common problem with older tool steels that reduces toughness and results in brittleness. Indeed, the presence of vanadium in conventional tool steels must be limited because vanadium tends to aggravate segregation. But vanadium carbide is particularly desirable for its high hardness and wear resistance, and new high alloy P/M grades have become available that form primary carbide levels (including vanadium carbide) exceeding 25% metal carbides by volume. Thus, the P/M grades of tool steels approach the hardness and wear resistance of cemented carbides while minimizing or eliminating many of the problems associated with cemented carbides.
A typical process of forming P/M grade tool steels comprises induction melting of a pre-alloyed tool steel composition, followed by gas atomization to produce a rapidly solidified spherical powder. This powder may then be applied to a base steel substrate by either weld overlay or, preferably, by hot isostatic pressure (HIP). Of course, the substrate could be eliminated if P/M powder were used to form an entire structure such as a valve seat by use of HIP (i.e., by HIPPING), but the cost of a valve seat comprising 100% of P/M grade tool steels would be prohibitive. And in spite of its high cost, such a valve seat would lack the toughness and strength otherwise obtainable if mild steel or a lower grade tool steel were used as a substrate.
HIP is a preferred method of applying a P/M grade inlay to a substrate because welding degrades some of the desirable properties of P/M tool steels. Even when welded ideally, the P/M grades will lose their fine microstructure in the weld fuse zone, where they melt during welding. Thus, P/M grades, when welded, do not achieve optimal toughness. Further, the melting that occurs during welding will decarburize some of the carbides, decreasing wear resistance. For these reasons, using welding to apply the high alloy P/M grades on heavy impact areas such as a valve seat will always present some risk of cracking in service. Rather, to make best use of high alloy P/M grades, they must be applied by HIP.
The HIP process avoids problems associated with welding because HIP is carried out at a temperature that is slightly lower than the melting temperature of the material being HIPPED. In fact, the ideal HIP temperature is the temperature at which the HIPPED material is only slightly plastic.
To economically produce a HIPPED valve seat with high alloy P/M tool steel grades, a suitable material such as H13 tool steel must be selected to serve as the seat substrate material. Traditional valve seat materials are not acceptable because of the high temperature required to harden the new high alloy P/M tool steel grades (i.e., about 1875° F. to 2200° F.). These temperatures are well above the austenitizing temperatures of traditional seat materials and can therefore lead to degraded grain structure in such substrate steels (which should only be hardened at a maximum of about 1750° F. to preserve desired grain structure). If a desired grain structure is destroyed by overheating, the affected steel reverts to a traditional cast grain structure with large coarse grains, causing loss of toughness and a tendency to crack under the cyclic fatigue impact loading typical of valve seats. Additionally, a structure comprising a tool steel substrate having a tool steel inlay can be quenched as a unit by inert gas in a vacuum furnace (thus preventing oxidation and/or decarburization of the tool steels). A substrate of traditional carbon steel, on the other hand, must be quenched in liquid (preferably oil) which would degrade a tool steel inlay.
But if a high alloy P/M tool steel grade were to be applied to a traditional steel substrate and hardened as necessary for tool steel, the steel base would not harden properly and would not develop the physical properties necessary to support the tool steel inlay under impact loading. Hence, a tool steel substrate such as H13 is preferable for the reasons given above.
Further, H13 is hardened at temperatures between 1600° F.–1925° F. These temperatures overlap with the hardening temperatures noted above for the new high alloy P/M tool steel grades. Thus an H13 seat, with a HIP-applied inlay of the new high alloy P/M tool steel grades can be hardened at 1900° F. and then tempered to obtain the ideal physical properties of both steels. This results in a seat with superior wear resistance on areas subject to high impact, as well as excellent strength and toughness in other areas of the seat. Areas that mate with the valve body's elastomeric seal and areas where the valve guide slides up and down during opening and closing of the valve also benefit from the improved wear resistance of the H13 substrate. These benefits are due in part to the fact that when an H13 substrate is hardened at 1900° F. it develops carbides throughout. Though such carbides are present at a much lower level than in the P/M grades of the inlay, they add substantially to the overall durability of the valve seat.
In current industry practice, HIP-applied inlays as described above require that the P/M powder be subjected to heat and pressure in a sealed enclosure (e.g., a metal can) which is evacuated to less than 0.1 torr (i.e., less than 0.1 mmHg). Empirical data show that this high vacuum is needed to reduce the inlay's porosity to achieve an inlay density of at least 99.7%. High density of the inlay is necessary to prevent formation of porous defects in the finished valve seat. Such porous defects, if present under cyclic fatigue impact loading, act as stress risers which lead to cracks, crack propagation, and catastrophic failure. Establishment of a high vacuum within the sealed HIP enclosure reduces these problems and also avoids undesirable oxidation of both the tool steel substrate and the P/M powder inlay during subsequent heat treatment.
In some pre-HIP applications, P/M powder may be preformed into a shape corresponding to the final inlay position on the substrate. This preforming is generally done independent of the substrate itself. Powder preforms are commonly made using a Cold Isostatic Pressure (CIP) process in which the powder is forged into a physical shape that, while porous (typically about 50% voids), is held intact at the inlay position by mechanical bonds among the powder particles. Typically, CIP is applied by placing the P/M powder in some type of deformable mold (e.g., rubber) having the desired shape and then pressurizing the mold. The pressurized deformable mold then collapses on the powder, compressing it under very high pressure (typically at least 30,000 psi.). Higher grade tool steel P/M powders generally require relatively high CIP pressures to achieve the necessary structural integrity for a powder preform to prepare it for subsequent application of HIP. This is because the greater hardness of these tool steel P/M powders makes the powder particles relatively resistant to the deformation required to achieve sufficiently strong mechanical bonds among the particles. After the preform is subjected to CIP, it is then assembled or mated with the corresponding substrate part inside a deformable metal can in preparation for the HIP process.
A metal can used for application of HIP may, if it provides complete sealing around a powder inlay and substrate, facilitates evacuation of the space adjacent to the inlay as described above. For example, the can used in current industry practice for the powder preform has welded seams and completely surrounds both inlay and substrate. A cross-section of such a typical welded can assembly with its enclosed valve seat substrate and inlay is shown in FIG. 4.
Note that a welded can assembly analogous to that of FIG. 4 usually has an evacuation tube. When present, such a tube allows evacuation of the can assembly after it is welded together (with the evacuation tube then being crimped/welded shut to maintain the vacuum within the can assembly). If a can assembly does not have an evacuation tube, this means that the can assembly itself must be welded together in a high-vacuum environment using a technique, such as electron beam welding, which is suitable for welding in a vacuum.
Since the can assembly in FIG. 4 does have an evacuation tube, it may be welded together using conventional techniques. The welded can assembly is tested for leaks with helium, after which the helium and any residual air are then evacuated via the evacuation tube. After evacuation, the evacuation tube is first crimped shut and then welded. The evacuated can assembly is then placed in a HIP furnace that is pressurized (typically with an inert gas) to a pressure of about 15,000 psi. Simultaneously, induction coils inside the HIP furnace heat the evacuated can assembly to a temperature just below the melting point of the parts, typically about 2200° F. for tool steels. The pressurized evacuated can assembly is held at this temperature for approximately four hours, after which the P/M tool steel powder has been solidified and forged into an inlay having a metallurgical bond (i.e., fused) with the tool steel valve seat substrate.
Note that the currently practices of various versions of the basic CIP process described above are all relatively expensive. High costs are associated with the molds and the tooling for the upper and lower portions of the can assembly, as well as the special handling required in welding, pressure testing, evacuating, crimping, and sealing can assemblies. In fact, the cost of preparing evacuated can assemblies as described above may substantially exceed the cost of applying HIP to these same assemblies.